Iron and steel materials containing Ni and Cr include Ni--Cr steel and stainless steels. As is well known, there are many types of stainless steels having superior corrosion resistance, weatherability, oxidation resistance, weldability, cold-workability, machinability and work-hardenability. These materials are used extensively in various chemical industries, architecture, turbines and related structures, aircraft, vehicles, etc. The stainless steels fall within four groups: austenitic, ferritic, martensitic and precipitation-hardening. Each group has its own advantages and disadvantages. For example, martensitic stainless steel has high strength and hardness. But since its Cr content is as low as about 13 atomic % or its carbon content is as high as about 3 atomic %, it has inferior corrosion resistance to austenitic and ferritic stainless steels, and also has inferior fabricability in deep drawing, cold-forging, etc. Austenitic stainless steel, in spite of its superior corrosion resistance, has a tensile strength of as low as about 60 kg/mm.sup.2, and even when it is work- hardened, its strength is not great.
To increase toughness and workability, the finely divided crystal grains are treated. Unlike ordinary steels, crystal grains of austenitic stainless steel are not easily divided finely by heat-treatment. Accordingly, by hot-working, crystal grains of its fabricated articles tend to become extremely coarse, which is not desirable.
Ferritic stainless steel is less expensive than austenitic stainless steel, but is disadvantageous in respect of workability or corrosion resistance.
HH and HK steels in accordance with the standards or ACI (Alloy Casting Institute) are known as materials having increased high-temperature strength which are obtained by increasing the C content of Ni--Cr type austenitic steel. These steels do not appreciably raise a hot-workability problem. But since they are usually converted into products by casting, the productivity is low. Furthermore, because they contain a large amount of C and coarse carbide particles, they have substantially inferior creep strength and fatigue life time in thermal environment to SUS 347, etc.
Piano wires and maraging steel are metallic materials showing high tensile strength. However, since they contain coarsened carbide and precipitate particles, hot- and cold-working steps for imparting work-hardening, etc., become complex. In particular, ultrafine wires, when stretched, lack ductility, and tend to snap on elongation.
A method of producing continuous fine steel wires is disclosed in Japanese Patent Publication No. 39338/1979 (U.S. Pat. Nos. 3,861,452 and 3,933,441). These patent documents are directed to an Fe--C--Si--Mn--O alloy, and the ranges of the suitable amounts of Si and Mn are limited in order to solidify the molten state alloy in a cooling medium. For example, it is shown that steel wires could be formed from several types of examples including 97.7Fe-0.7Si-0.4Mn-1.2C and 93.5Fe-2.3Si-1.2Mn-3C. The patent documents state that since a solid silica precipitate forms upon contacting of the jetted molten steel with the cooling medium, the oxidation product acts as a hardening initiator and a hardening promoter. It is clearly described there that to obtain continuous fine steel wires, the preferred amount of Si is about 1 to 6 atomic %, and the preferred amount of Mn is about 0 to about 1.5 atomic %. These patent documents are quite silent on the properties of the steel wires produced, and merely describe alloy compositions which can give such fine steel wires.
Japanese Patent Publication (unexamined) No. 3651/1981 (Metall. Trans. A., Vol. 12A, p. 1245 (1981)) discloses imparting of toughness to an L1.sub.2 -type intermetallic compound. Specifically, it discloses an intermetallic compound material which is composed of 3.9 to 67.0 atomic % of at least one of Ni and Mn, 7.2 to 22.5 atomic % of Al, 0.7 to 11.0 atomic % of C or 0.7 to 11.0 atomic % of C and N (not more than 0.8 atomic %), and the balance being Fe and is almost entirely made up of an L1.sub.2 -type intermetallic compound and in which most of C or both C and N are dissolved in the intermetallic compound. It further states that not more than 7.4 atomic % of at least one of Cr, Mo and W may be added to the above alloy and not more than 42.0 atomic % of Ni and Mn may be replaced by Co. It is further stated that not more than 3.8 atomic % of at least one of Ti, Ta, Zr, Nb and Si may be added to the alloy, and even when Cr, MO, W, Co, Ti, Ta, Zr, Nb and Si are added, the alloy material is almost entirely composed of an L1.sub.2 -type intermetallic compound, and most of C or C and N are dissolved in the intermetallic compound. Because of its low Cr content (not more than 7.4 atomic %), high Al content and high C content, this intermetallic compound material has an ordered structure and an inverse phase area and exhibits toughness. This intermetallic compound material has toughness only within the aforesaid composition range. When the amount of Al is less than 7.2 atomic %, no L1.sub.2 -type intermetallic compound is formed, and the resulting material has low strength. When it exceeds 22.5 atomic %, an L1.sub.2 -type intermetallic compound is formed. But its toughness is markedly reduced and it becomes brittle. When the amount of Ni is less than 3.9 atomic %, the resulting material has markedly reduced toughness owing to the formation of a carbide. When it exceeds 65.5 atomic %, Fe.sub.3 C is formed and the resulting material loses toughness. When the C content is less than 0.7 atomic %, no appreciable quenching effect is observed. Hence, an L1.sub.2 -type intermetallic compound cannot be formed and the resulting material becomes brittle. When the C content is more than 11 atomic %, the precipitation of Fe.sub.3 C is difficult to prevent even by quenching, and the resulting material remarkably loses toughness and becomes brittle. Thus, the L1.sub.2 intermetallic compound material has toughness only within the aforesaid composition range, and outside this range, the precipitation of a carbide occurs, and the resulting material completely loses toughness and becomes brittle so that it is useless in practical applications. The L1.sub.2 -type intermetallic compound material having the above alloy composition has toughness, but is difficult to work by wire drawing, rolling, heat-treatment, etc. Moreover, an improvement in mechanical properties, etc., by working can hardly be expected. For example, an Fe.sub.59.8 Ni.sub.16.4 Al.sub.14.2 C.sub.9.6 alloy material having a tensile strength of about 175 kg/mm.sup.2 which is the highest among the aforesaid L1.sub.2 -type intermetallic compound materials contains many anti-phase boundaries and has fine anti-phase domains. Hence, even when it is subjected to some after-treatment without work-hardening, its tensile strength and yield strength cannot be improved over those of a quenched material. Furthermore, this intermetallic compound material is a non-equilibrium phase material. Hence, when it is heat-treated, for example, at 600.degree. C. for 1 hour, anti-phase boundaries abruptly vanish and consequently anti-phase domains essential to toughness vanish. As a result, it becomes an L1.sub.2 -type intermetallic compound of an equilibrium phase which loses toughness, becomes brittle, and is considerably unstable thermally. In addition, the intermetallic compound material has considerably low corrosion resistance since it has a kind of boundary called an anti-phase boundary in the grains and is an ultrahigh carbon material.